The deposition of heteroepitaxial Ge and Ge.sub.x Si.sub.1-x alloys on Si has attracted widespread interest because of the possibility of tailoring the band gap of Ge.sub.x Si.sub.1-x based heterostructures. Proposed uses for these heterostructures include optical devices, modulation doped transistors, and heterojunction bipolar transistors. For example, thin films of elemental Ge grown on Si substrates may be utilized as an intermediate epilayer to allow subsequent growth of gallium arsenide (GaAs) for optoelectronic integration on Si or for use as a SiGe heterojunction in high performance electronic device applications. Since GaAs is better lattice matched to Ge and can be easily deposited on it, a high quality epitaxial Ge layer on Si serves as a suitable substrate for the subsequent growth of GaAs and possibly other layers (e.g. Ga.sub.x Al.sub.1-x As), and consequently allows for the integration of GaAs optoelectronic devices with Si based devices.
Heteroepitaxial growth of pure Ge is most desirable, in particular, because the higher the Ge content of the structure, the higher the carrier mobility which, in turn, provides higher current capability, lower power consumption, lower temperature operation and greater speed. Also, as previously mentioned, pure Ge provides optoelectronic capability.
A complicating factor in fabricating Ge-Si heterojunction systems, however, is a large mismatch in both lattice parameter (.congruent.4%), and thermal expansion coefficient (.congruent.50%). Previous attempts to produce Ge-Si heterostructures exhibit a marked variation in the growth morphology and crystalline quality attributable to contamination presence at the Ge-Si interface, which is determined primarily by the initial condition and preparation of the Si substrate surface and improper deposition process parameters.
Ge and Ge.sub.x Si.sub.1-x alloys have previously been deposited on Si by a variety of techniques, including ion-sputtering, physical vapor deposition, chemical vapor deposition (CVD), and more recently, molecular beam epitaxy (MBE). CVD is the most advantageous of these processes because it is a high throughput process and also because it has in-situ doping capabilities. However, the high temperatures, high pressures and carrier gases characteristic of conventional CVD systems produce undesirable surface roughness and therefore render the process impractical for many applications.
Examples of prior methods of heteroepitaxial growth of Ge on Si utilizing CVD are disclosed in the following publications.
Maenpaa, et al., "The heteroepitaxy of Ge on Si: A comparison of chemical vapor and vacuum deposited layers", J. Appl. Phys. 53(2), (February 1982), 1076-1083, (Maenpaa et al.) describes a CVD experiment performed utilizing GeH.sub.4 as a carrier gas with a pressure of 2-13 Torr and maintaining the Si temperature between 500.degree.-900.degree. C.
Hitoshi Ishii, "Manufacture of Semiconductor Device and Equipment Therefor", Japanese Patent publication 62-179113, application no. 61-20283, (Ishii) discloses the deposition of Ge using GeH.sub.4 as a carrier gas in a CVD chamber at 450.degree. C.
U.S. Pat. No. 3,473,978 (Jackson), entitled "Epitaxial Growth of Germanium", teaches a method for the nucleation and growth of monocrystalline germanium on a Si substrate which comprises epitaxially growing a layer of monocrystalline Si at a temperature of at least 900.degree. C., then cooling the Si below 670.degree. C., followed by the nucleation and growth of Ge.
More recently, ultrahigh vacuum chemical vapor deposition (UHV/CVD) processes have been utilized for growing of Ge.sub.x Si.sub.1-x layers onto Si. Examples of this technique are disclosed in the following references.
Meyerson, Uram and LeGoues, "Cooperative growth phenomena in silicon/germanium low-temperature epitaxy", (Appl. Phys. Lett. 53 (25), Dec. 19, 1988, 1988 American Institute of Physics), (MEYERSON ET AL.), which teaches deposition of alloys of composition 0.ltoreq.Ge/Si.ltoreq.0.20 using UHV/CVD and a temperature of 550.degree. C.
Racanelli and Greve, "Growth of Epitaxial Layers of Ge.sub.x Si.sub.1-x by UHV/CVD", (Mat. Res. Soc. Symp. Proc. Vol. 198, 1990 Materials Research Society), (RACANELLI et al.), which teaches a method of growing epitaxial layers of a Ge.sub.x Si.sub.1-x composition on Si at temperatures between 577 and 665.degree. C.
In CVD and UHV/CVD processes of the prior art, a reactant gas is transported over a heated substrate. A chemical decomposition reaction occurs on the substrate and deposition occurs with a growth rate which is variably dependent upon the substrate temperature, wherein two distinct temperature (high and low) regimes are found to exist. In the low temperature regime, deposition is strongly dependent upon temperature, while only a small dependence is shown for the high temperature regime. FIG. 1 illustrates a typical CVD growth rate versus temperature curve. For the low temperature regime, wherein deposition is kinetically controlled via a surface decomposition, the overall growth rate G follows an Arrhenius plot according to: EQU G=G.sub.o exp (-.DELTA.E/kT.sub.s)
where:
G.sub.o is a pre-exponential factor,
.DELTA.E is the apparent activation energy PA1 k is Boltzmann constant PA1 T.sub.s is the substrate temperature.
The value of the activation energy can be used to postulate the rate limiting step of the surface reaction or, conversely, if the rate limiting step is known, to determine the activation energy for such a reaction.
It is to be noted that above the transition temperature, T*, the growth rate is mass transport/diffusion-limited. In this regime the rate limiting step is the diffusion of the gaseous reactant to the surface, since the high deposition temperature rapidly decomposes and incorporates the molecules arriving on the surface. This fast decomposition step makes available a sufficient number of surface sites to allow for the continuous adsorption of the arriving reactant species. Since the diffusion coefficient varies slowly with temperature, a slower increase in growth rate is observed in the diffusion limited regime as compared to the kinetically controlled regime. This behavior has been observed for CVD of epitaxial Si from a variety of gas sources (e.g. SiH.sub.x Cl.sub.1-x).
As previously mentioned, it is the heteroepitaxial growth of Ge which is desirable. FIGS. 2 A-C show schematic representations of the three basic modes of heteroepitaxial growth. The first mode (FIG. 2A) is known as Frank-van der Merwe (F-vdM) wherein growth proceeds in a layer by layer mode (2D). The second mode (FIG. 2B) is known as Volmer-Weber (V-W), wherein island growth (3D) occurs. The third mode (FIG. 2C) is known as Stranski-Krastanov (S-K), which proceeds initially as layer by layer growth, followed by islanding after a critical thickness has been exceeded. In the absence of any strain in the epilayer (i.e. for lattice matched systems), the growth mode is determined by the surface energy of the substrate (.sigma..sub.1) and epilayer (.sigma..sub.2), and the interface energy (.gamma..sub.12) and is predicted purely by the following thermodynamic free energy considerations: EQU .sigma..sub.1 &gt;.sigma..sub.2 +.gamma..sub.12 F-vdM (2D) EQU .sigma..sub.1 &lt;.sigma..sub.2 +.gamma..sub.12 V-W (3D)
When the epilayer has a large mismatch to the substrate, as is the case for Ge on Si, an additional term must be included to account for the strain energy of the epilayer. This strain energy (.epsilon..sub.t) increases as the epilayer thickness increases, and the expressions used to determine the growth mode then become: EQU .sigma..sub.1 &gt;.sigma..sub.2 +.gamma..sub.12 +.epsilon..sub.t F-vdM(2D) EQU .sigma..sub.1 &lt;.sigma..sub.2 +.gamma..sub.12 +.epsilon..sub.t V-W(3D)
Since .epsilon..sub.t is positive, systems which are expected to grow in a V-W mode in the absence of strain will not be affected by the additional strain energy term. It is, however, possible for the system to begin growth in the F-vdM mode and change to 3D growth if the strain energy exceeds a critical value. This mode of growth is the S-K mode, and is shown schematically in FIG. 2C.
It should be noted that the preceding theoretical model of the growth modes is only valid in a strict sense, that is, if no chemical reactions or other changes on the surface occur during film growth. In other words, the predicted growth mechanisms are based purely on energy considerations.
The surface free energy of Ge is smaller than that of Si by about 10%, which would imply that the growth of Ge/Si should proceed in a F-vdM mode, However, since the misfit between the two lattices is approximately 4%, the accumulation of strain energy, in fact, produces S-K growth, with 3-4 monolayers (ML) of layer by layer growth occurring prior to island formation. Prior techniques, (such as those mentioned hereinbefore), for depositing Ge on Si have been undesirable because these factors have produced rough surface morphology, as illustrated in FIG. 3.
In addition to the production of films having rough surface morphology, CVD methods of the prior art are undesirable because they require high processing temperatures, thereby limiting device application. Also, prior CVD methods utilize poor vacuum processing techniques, which cause system contamination, thereby affecting surface interface quality essential to film growth.
Likewise, none of the aforementioned UHV/CVD references describe a suitable method for growing pure Ge on Si.
It is possible, however, to amend the aforementioned Ge growth mechanisms by changing the equilibrium conditions of the surface during growth. A system which accomplishes this and overcomes the deficiencies of the prior art is highly desirable.